The covalent nature of the silicon-carbon bond imparts the intrinsic source of the mechanical strength and stiffness of SiC articles, and accounts for the low self-diffusivity of silicon and carbon atoms in SiC over a very wide temperature range. This is especially the case for high purity CVD SiC. This property also plays a dominant role in establishing the difficulty with which the SiC can be bonded to itself as well as to other ceramics and metals.
The CVD process can produce a condensed form of SiC via the decomposition of as phase reactants at comparatively low temperatures, e.g., 1100° C.-1400° C., without the use of any sintering or other densification aids. This implies that high purity SiC can be produced by this synthesis route. Indeed CVD SiC is a particularly desirable form of SiC, producing articles having properties including, but not limited to, essentially zero porosity, extremely high purity (>99.9995%), no need for second phases or sintering aids, and a very wide range of possible resultant component shapes.
Typically, methyltrichlorosilane (CH—SiCl3), or ethyltrichlorosilane (C2H5SiCl3) gas, employed with hydrogen as a carrier gas, is employed as a gas phase route for the synthesis of bulk, often very high purity, CVD SiC articles. The CVD SiC is typically synthesized by gas phase decomposition of gaseous precursors according to the following reaction paths:CH3SiCl3+H2→SiC+3HCl (for MTS)C2H5SiCl3+H2→SiC+3HCl+CH4 (for ETS)
Deposition rates are generally fairly slow (on the order of 0.0025 cm per hour), and crystalline microstructures tend to be columnar in morphology, with the grains being smaller at the onset of deposition and then growing larger as preferred crystalline variants begin to grow in the thickness direction. Given the deposition temperatures. CVD SiC is essentially composed entirely of β-phase material.
A phase transformation from β-SiC to α-SiC occurs upon heating, with undoped β-SiC transforming to 6H and 15R in the vicinity of 2000° C., and the 15R polytype being a metastable phase that transforms to 6H. Doping β-SiC with boron lowers the transformation temperature and promotes formation of the 4H polytype. On the other hand, doping with nitrogen prevents the formation of 4H and stabilizes the 6H. The β→α phase transformation is irreversible at ambient pressures. However, under pure nitrogen atmosphere, the transformation can be reversed and β-SiC phase can be stabilized up to 2500° C. by applying processing in nitrogen ambient atmospheres. At atmospheric pressure and temperatures exceeding 2000° C., pure silicon, carbide does not melt, but sublimes or dissociates.
Referring to FIG. 1, a representation of α-SiC (a) vs. β-SiC Crystal Structure (b) is shown. The density if SiC for all polytypes is 3.21 g/cc and SiC has a melting (actually decomposition) point of 2730° C. As a general matter, CVD SiC has several desirable engineering attributes including, but not limited to: CVD SiC has high hardness in the range of 2200-2500 kg/mm2; CVD SiC has a high elastic modulus of 400-460 GPa; CVD SiC possesses excellent dimensional stability and stiffness as well as outstanding polishing characteristics (e.g., polishability to <3 Å RMS); CVD SiC has a low coefficient of thermal expansion (CTE) (e.g., 2.2-4.5×10−6/C from RT to 1500° C.); CVD SiC has excellent elevated temperature strength retention characteristics (e.g., flexural strengths of typically 400-420 MPa at room temperature, while peaking at strengths of 550-580 MPa at temperatures in the vicinity 1400° C.); and CVD SiC has excellent thermal conductivity (e.g., 100-300 W/m-K including for certain doped grades (e., Dow Chemical CVD SiC®)).
In part for these reasons, CVD SiC has filled important, growing product niches in demanding technical applications, which include tactical and spaceborne IR/optical mirrors, laser galvanometers semiconductor wafer carrier components and specialty heat sink components. Applications for the bonded SiC articles of the present invention can be anticipated in the semiconductor, optics, optoelectronics, petrochemical, nuclear energy realms, and the like.
In spite of its favorable overall mix of technical properties CVD SiC has some engineering and practical shortcomings, which include but are not limited to: CVD SiC is difficult to make into bonded structures (particularly to itself), due to its low lattice and boundary self-diffusion coefficients (Dlattice˜2.62±1.83×108 {−8.72±0.14 ev/atom/kT}cm2 sec−1; Dboundary˜4.44±2.03×107 {−5.84±0.09 ev/atom/kT}cm2 sec−1). See, for example, Hon, M. H., and Davis, R. F., “Self-Diffusion of 14C in Polycrystalline —SiC,” Journal of Materials Science, 14 (1979) 241-2421. It is possible to bond CVD SiC members to themselves by making use of methods which may employ refractory metallic foils (i.e., Zr, Ti, Mo, etc) or pre-ceramic silazane-type polymers, but such methods will impose significant limitations on the temperature capability of the bonded structures due to formation low-melting point eutectic phases such as TiSi2 (melting point: 1470° C.) or otherwise due to the intrinsic performance limitations (decomposition, loss of elevated temperature strength, etc.) of the joint material itself.
Conventionally grown CVD SiC has thickness limitations (e.g., ˜1-2 cm) associated with the accumulation of residual stress for thick deposits. It is in fact sometimes observed that deposits of over 1.2 cm-1.3 cm thickness can spontaneously fracture due to buildup of crystalline growth stress, greatly complicating the machining and fabrication of thick articles. This is an unfortunate aspect of CVD SiC fabrication, as there is demand for thick articles having complex geometrical features for applications such as the semiconductor and lightweight optics industries.
Current methods, as in U.S. Pat. No. 4,925,608, for example, teach joining sintered α-phase silicon carbide via a hot isostatic press (HIP) diffusion banding method. The method entails polishing surfaces to a mirror finish with diamond abrasives, fitting the surfaces together to form a composite structure, and then subjecting the composite structure to HIP processing (pressure-assisted diffusion bonding) under conditions that promote plastic deformation and diffusion flow at the bonding interface. While capable of producing well bonded and geometrically consistent metallic encapsulation layers, this method is costly and limited with regard to its ability to accommodate complex shapes. Difficulties with transition to huge-scale manufacturing is also seen, as expensive restraint tooling and access to HIP furnaces is required, thus limiting applications of this method to simple geometries in applications that are not cost-sensitive.
Another current method, as in U.S. Pat. No. 5,683,028, teaches joining CVD SiC via the use of geometrically profiled male and female joint members enclosing a gap, which is approximately 0.076 cm in width. There, the female joint also has a reservoir provided which accepts a volume of molten Si during infiltration and wicking of liquid Si into the joint region during thermal processing at temperatures of approximately 1410° C. (the melting point of Si). Following liquid Si infiltration, the region joined by molten Si is allowed to solidify and is then overcoated with a layer of CVD SiC to isolate the silicon-bonded region at the interface. This method, while capable of providing a joint with high strength, is procedurally complex, costly, and would impose thermal limitations on the joined parts, given that the melting point of silicon is 1410° C. Thus, in cases where use of the bonded article is required above the melting point of Si, the method would cause excessive weakening of the joint due to plastic deformation and liquid phase formation. Additionally, excessive oxidation at temperatures over 1300° C. would result due to the presence of Si-rich regions in the joint. This method would also not be particularly suitable for use in a field-expedient setting such as might be required for repairing or joining subassemblies.
Additionally, U.S. Pat. No. 4,961,529 teaches a method for joining SiC (type not specified) in a fashion somewhat similar to U.S. Pat. No. 5,683,028, making use of geometrically profiled male and female joint members, but with an imposed layer of titanium carbosilicide (Ti3SiC2). There, a thin layer of Ti3SiC2 is made by sputtering or by mixing powders of TiC0.8 and TiSi2, and the layer is interposed between the profile SiC articles to be joined, which are pre-machined and highly polished. In that method, the actual bonding process occurs by vacuum hot pressing the articles to be joined at temperatures in the 1450° C.-1500° C. range. This method is procedurally complex and costly, while also being limited to joining articles that can be accommodated within a hot press die set, thus severely hampering practical applications to complex or large assemblies. Additionally, excessive oxidation and plastic deformation at temperatures over 1200° C. would also be likely due to the presence of the titanium carbosilicide bond phase. Additionally, the Ti3SiC2 phase has a significantly different elastic modulus and coefficient of thermal expansion (CTE) as compared to pure SiC, which would exacerbate accumulation of undesirable elastic and plastic strains in the joined articles—particularly in a cyclic stress or temperature environment. Just as the methods previously discussed, the techniques described in U.S. Pat. No. 4,961,529 would not be compatible with use in a field-expedient setting such as might be required for repairing or joining SiC subassemblies.
It is recognized that major deficiencies remain with state of the art methods used to bond SiC-based ceramics, and especially CVD SiC. Accordingly, there is a need to develop a method to bond CVD SiC in a manner that is less complicated, less costly, capable of working with a wide range of component sizes and geometries, and is also compatible with the requirements of reproducible and large-scale manufacturing.